High-carbon precipitation-hardening austenitic steel alloy



United States Patent 3,309,242 HIGH-CAREON PRECIPTTATION-HARDENING AUSTENITIC STEEL ALLOY Horace Pops, Wiikinsburg Borough, Pa, assignor to United States Steel Corporation, a corporation of Delaware No Drawing. Filed Oct. 1, 1964, Ser. No. 400,907 4 Claims. (Cl. 14838) This invention relates to high-carbon precipitationhardening austenitic iron alloys characterized by high temperature tensile strength suitable for aircraft and missile uses.

Aircraft and missile structures require alloys that will retain strength in the range of 1000 to 1500 F. Presently used materials generally are low-carbon, chromium, nickel, iron alloys containing titanium as a major hardening agent together With molybdenum, vanadium and alu minum. Strengthening of these alloys is accomplished by precipitation or age-hardening and is usually attributed to internal lattice strains produced by the precipitation of intermetallic compounds such as Ni (TiAl). At high temperatures the precipitates agglomerate to larger particles and thus lose their effectiveness. The carbon content of titanium-bearing high-temperature alloys is kept at a minimum since With increased carbon contents, the titanium acts as a carbide stabilizer and forms relatively stable titanium carbide dispersoids so that the matrix is depleted of titanium and is therefore not available as a precipitation-hardening element. High-carbon, high-temperature iron alloys are also known; these are usually strengthened by the precipitation of chromium carbides and are used mainly for valves and the like. However such steels have inferior mechanical properties at the higher temperatures associated with aircraft and missile applications.

It is an object of this invention to provide an austenitic iron alloy having a considerably higher level. of strength for high temperature service because the zirconium carbide particles, which form, become agglomerated and segregate in the as-cast ingots. On the other hand, -I have discovered that with relatively high amounts of carbon and zirconium, a large number of uniformly dispersed zirconium carbide dispersoids are formed and consequently no internal stress raisers are produced because of the uniform distribution. In addition to this dispersion strengthening mechanism, one or more phases precipitate out of the matrix during aging heat treatments thereby extending the high strength-characteristic of the steel over longer periods at elevated temperature. Thus the mechanical properties of the alloy of this invention, are believed to be obtained by a superimposition of multiple strengthening mechanisms. Moreover despite the large volume fraction of carbide particles, the alloy of this invention can be readily hot or cold swaged or machined.

The austenitic iron alloy of my invention has the following compositional range by weight'precent:

purities. In the latter regard, residuals and impurities include not only the normal small amounts of phosphorus, sulphur and dissolved gases present in steels but also manganese, silicon and aluminum commonly presentv as the results of manufacturing processes; the latter elements being unnecessary to the properties of the present alloy but not deleterious thereto in amounts up to 1.0% each.

The following heats of steel will illustrate the invention:

TABLE I ltd Bal. Bnl. Bal. Ba].

coww

mmtn

and improved ductility at temperatures above about 1200 F. than available heretofore.

A further object of this invention is to provide a steel having a matrix that is strengthenable by dispersion of relatively stable carbide particles as Well as by precipitation hardening.

I have discovered that the foregoing objects can be obtained in an austenitic chromium nickel steel containing high carbon within a restricted range together with limited amounts of zirconium, titanium, molybdenum, tungsten and cobalt; further that the inclusion of a small amount of boron in the composition is very beneficial. Small additions of zirconium are not customary in iron alloys The tensile strength at various test temperatures of the foregoing steels after precipitation hardening is listed in the following Table 11 along with that of A-286- steel which is sometimes used for high temperature service. A-286 is a commercially available steel of the loW carbon, chromium, nickel, titanium and iron type. As apparent upon comparison of the properties of steels 4 and A-286, the present invention afifords materials of considerably higher strength at elevated temperature than avaiiable heretofore. The ultimate strength of steel 4 is about greater at 1500 F. than that of A286 while its yield strength at this temperature is about greater. Moreover steel 3 Which is essentially the same as steel 4, except for the presence of 0.004% boron affords an even more remarkable result showing improvements in both tensile and yield strengths of better than 45% at 1300 F. and of more than 100% at 1500 F. As is indicated by comparing the elongations of steels 3 and 4, boron markedly increases the elevated temperature strength of the alloys without lowering the ductility. The effects of boron however are lost when the carbon content of the alloys is increased, see steel 1. That the carbon content must be closely controlled is further evident from steels 2 and 3, wherein a relatively small increase from .57 to .68% has resulted in a large drop in the tensile strength.

TABLE II.-TENSILE AND YIELD STRENGTHS AND ELONGATION FOR STEELS 1, 2, 3, 4 AND A-ZSG AT 'IEhi- PERATURES INDICATED Test Tensile Yield Average Steet Tempera- Strength Strength Elongation ture (l .)j (p.s.i.) 02% Ofiset in 1",

(p.s.i.) Percent 1 1 68 143, 700 92, 000 10.3 1, 000 113, 300 84, 600 11.0 1, 300 104, 400 91, 600 9.6 1, 500 47, 000 44, 400 15.7 1, 650 17, 500 16, 800 36.5 2 1 68 140, 000 111, 100 13.4 1, 000 114, 800 90,100 9.2 1, 300 108,100 84, 300 11.5 1, 500 37, 300 35, 100 18.1 1, 650 21, 000 20, 100 42.0 3 1 6B 183, 000 128,300 6.0 1,100 160, 900 124, 000 7.2 1, 300 131, 300 109, 700 9.4 1, 500 65, 600 57, 600 28.0 1, 650 21, 500 19, 000 34. 4 1 68 1 16, 000 08, 000 7.8 1, 000 125, 500 101, 400 9.2 1, 300 107,200 81, 300 13.6 1, 500 48, 100 43, 700 25.2 1, 650 20, 200 18, 900 30.0 A-ZSG Z G8 152, 000 98, 000 20.0 1, 000 130, 000 S7, 000 16. 0 1, 100 121, 000 S5, 000 13.0 1, 300 88, 000 75, 000 4. U 1, 500 31, 000 17,000 14.0 1, 650 18, 000 52.0

1 hr. at 1,850" F. 10 hrs. 1,300 F. 9 hr. at 1,800 F. 13 hrs. 1,326 F.

The composition of the A-286 steel included in the foregoing table .had the following composition:

The properties of the alloys of the present invention are developed by heating treatment involving a solution anneal at temperature ranging from about 1800 to 2200 F. followed by a precipitation or aging treatment at a lower temperature.

The effects of variations in treating conditions are illustrated by the data of Tables III, IV and V. While the tabulated data are specific to steel 3 when tested at 1300 F., they may be considered typical of all the compositions of the present invention as regards the variation in properties to be expected with variations in temperature of solution anneal and the time and temperature used in the precipitation treatment.

TABLE IIL-EFFECT OF AGING TIME UPON HIGH TEM- PERATURE TENSILE STRENGTH AT 1,300 F. FOR STEEL 3 Solution Aging Treatment Treatment Yield Average Tensile Strength Elong. Strength 02% in 1, Temp. Time Temp. Time (p.s.i.) Oilset percent F.) (hrs) F.) (hrs) (p.s.i.)

1, 850 1, 300 1 120, 300 93, 800 10.8 1, 850 1, 300 10 131, 300 109, 700 9.4 1, 850 V 1, 300 133, 900 106, 500 14. 7 1, 850 A 1, 300 400 118, 800 95, 600 12.1 1, 850 1% 1, 300 750 114, 200 92, 300 13.6 1,850 1, 300 1, 000 113, 200 86, 900 14.2 2, 000 L6 1, 300 1 111, 600 85, 500 10.0 2, 000 A 1, 300 10 124, 700 113, 800 9 0 2, 000 1, 300 100 127, 800 113, 500 1'3 7 2, 000 1g 1, 300 400 93, 600 89, 000 15. 9 2, 000 A 1, 300 750 108, 100 87, 000 22.4 2, 000 i l, 300 1, 000 91, 200 75,100 14.0 2, /2 l, 300 0 02, 500 52, 100 22.3 2, 150 1, 300 1 89, 600 53, 800 12.1 2, 150 3 1, 300 10 106,100 83, 300 11.7 2, 150 1, 300 100 117, 900 9d, 100 13.9 2, 150 1,300 400 118,100 101, 000 15.0 2, 150 A 1, 300 750 113, 300 92, 300 28.5 2, 150 1, 300 1,000 94, 800 81, 100 13.2

TABLE IV.EFFEGT OF AGING TEMPERATURE UPON HIGH TEMPERATURE TENSILE STRENGTH AT 1,300" F. FOR STEEL 3 Solution Aging Treatment Treatment Yield Average Tensile Strength Elong. Strength 02% in 1", Temp Time Temp. Time (p.s.1.) Ofiset percent F.) (hrs) F.) (hrs) (p.s.i.)

1,850 )6 l, 200 10 122, 700 94, 000 15.9 1, 850 1, 300 10 131, 300 109, 700 0.4 1, 850 l, 400 10 124, 500 98, 700 7.9 1,850 A l, 500 10 95, 000 69, 400 14. 5 1, 850 M), 1, 600 10 94, 200 69, 000 22 1 2, 000 l, 200 10 90, 600 66, 000 15.5 2, 000 l 1, 300 10 124, 700 113, 800 9.6 2, 000 1, 400 10 93, 700 80, 300 9.1 2, 000 A 1, 500 10 71, 700 54, 400 17.6 2, 000 1, 000 10 71, 400 43, 900 19.1 2, 150 1, 200 10 99, 500 66, 300 23 2, 150 1, 800 10 106, 100 83, 300 11 7 150 1, 400 10 111, 800 87, 100 6.8 2, 150 1, 500 10 90, 500 700 11.8 2, 150 l, 600 10 98, 500 60, 400 22.4

In general and as evident from Table III, solution annealing at temperatures toward the lower end of the range 1800 to 2200 F. will produce somewhat higher levels of strength. Annealing time is preferably held to the minimum needed to achieve substantially uniform temperature throughout the alloy. While this will vary with thiclo ness of section, a half hour at temperature is in most instances sulficient.

As also evident from Table III, considerably longer times are involved in the precipitation treatment; aging for 100 hours or more at 1300" F. being required to develop maximum strength. However it will be noted that the increase in strength with times longer than 10 hours is minor; consequently treatment for longer than about 10 hours are uneconomical and may even be detrimental.

The effect of aging temperature at a constant treating time of 10 hours is illustrated in Table IV. It will be noted that the optimum aging temperature rises to about 1400 F. as the temperature used in the solution anneal is increased. More important however is the fact that the highest strength level is achieved by solution annealing at about 1850 F., in combination with aging at about 1300 F. It is apparent that lower aging temperatures, e.g. 1200" F., are usable but would require extension of the time of treatment, while higher aging temperature, e.g. 1500" F. should be avoided since the strength level drops even when the treating time is drastically shortened. The latter is illustrated by the data of Table V.

TABLE V.EFFECT OF TDAE AND TEMPERATURE CF AGING ON B/IECHANICAL PROPERTIES OF STEEL 3 AS DETERLHNED BY TENSILE TESTS AT 1,300 F.

Solution Aging Treatment Treatment Average Tensile Yield Elong.

Strength Strength in 1, Temp. Time Temp. Time (p.s.i.) (p.s.i.) percent F.) (hrs) F.) (his) I 1, 850 $6 1, 300 i 1 120, 300 93, 800 10.8 1,850 1,300 10 131,300 100, 700 9.4 l, 850 1, 500 1 103, 700 87, 500 10.8 1,850 1, 500 10 95. 000 69, 400 14.5 2, 150 l- 1, 300 1 80, 600 53, 800 12.1 2, 150 A 1,300 10 106, 100 83. 300 11.7 2, 150 1, 500 1 100,100 87, 600 14.8 2, 150 Li I 1, 500 10 90, 500 61, 700 I 11.8

In summary, optimum practices of the invention involve a short time solution anneal of the alloys disclosed at about 1850 F. followed by precipitation hardening by heating to about 1300 F. (12751350 F.) for about 10 hours.

When thus treated, the high-carbon precipitation-hardening austenitic steels of this invention are characterized by maintenance of high strength for more than 1000 hours at temperatures above 1200 F. and up to about 1500 F., a significant improvement over the presently available steels used for high temperature service.

While I have shown and described several specific embodiments of my invention, it will be understood that these embodiments are merely for the purpose of illustration and description and that various other forms may be devised within the scope of my invention, as defined in the appended claims.

I claim:

1. A precipitation-hardening austenitic iron alloy consisting essentially of, by weight percent Carbon 0.51 to 0.60 Chromium 10.0 to 20.0

Nickel 15.0 to 25.0 Zirconium 2.0 to 4.0 Titanium 2.0 to 4.0 Molybdenum 2.0 to 4.0 Tungsten 2.0 to 4.0 Cobalt 2.0 to 4.0

with the balance iron and incidental residuals.

2. A precipitation-hardening austenitic iron alloy consistin g essentially of, by weight percent Carbon 0.51 to 0.60 Chromium 10.0 to 20.0 Nickel 15.0 to 25.0 Zirconium 2.0 to 4.0 Titanium 2.0 to 4.0 Molybdenum 2.0 to 4.0 Tungsten 2.0 to 4.0 Cobalt 2.0 to 4.0 Boron Up to .008

with the balance iron and incidental residuals.

3. A precipitatiomhardened austenitic iron alloy consisting essentially of, by weight percent Carbon 0.51 to 0.60 Chromium 10.0 to 20.0 Nickel 15.0 to 25.0 Zirconium 2.0 to 4.0 Titanium 2.0 to 4.0 Molybdenum 2.0 to 4.0 Tungsten 2.0 to 4.0 Cobalt -Q 2.0 to 4.0

with the balance iron and incidental residuals.

4. A precipitation-hardened austenitic iron alloy consisting essentially of, by weight percent Carbon 0.51 to 0.60 Chromium 10.0 to 20.0 Nickel 15.0 to 25.0 Zirconium 2.0 to 4.0 Titanium 2.0 to 4.0 Molybdenum 2.0 to 4.0 Tungsten 2.0 to 4.0 Cobalt 2.0 to 4.0 Boron Up to .008

with the balance iron and incidental residuals.

References Cited by the Examiner UNITED STATES PATENTS 2,157,060 5/1939 Schafrneister -128 2,323,490 4/1945 Mohling 75128 2,617,725 11/1952 Owens 75128 2,677,610 5/1954 Evans 75l28 DAVID L. RECK, Primary Examiner.

P. WEINSTEIN, Assistant Examiner. 

2. A PRECIPITATION-HARDENING AUSTENITIC IRON ALLOY CONSISTING ESSENTIALLY OF, BY WEIGHT PERCENT 